Materials for electronic and optoelectronic devices having enhanced charge transfer

ABSTRACT

A composite material is described. The composite material comprises semiconductor nanocrystals, and organic molecules that passivate the surfaces of the semiconductor nanocrystals. One or more properties of the organic molecules facilitate the transfer of charge between the semiconductor nanocrystals. A semiconductor material is described that comprises p-type semiconductor material including semiconductor nanocrystals. At least one property of the semiconductor material results in a mobility of electrons in the semiconductor material being greater than or equal to a mobility of holes. A semiconductor material is described that comprises n-type semiconductor material including semiconductor nanocrystals. At least one property of the semiconductor material results in a mobility of holes in the semiconductor material being greater than or equal to a mobility of electrons.

RELATED APPLICATIONS

The present application is a continuation of U.S. patent applicationSer. No. 12/426,854, filed on Apr. 20, 2009 now abandoned, which claimsthe benefit of priority under 35 U.S.C. 119(e) to U.S. ProvisionalPatent Application Ser. No. 61/046,390, filed on Apr. 18, 2008, U.S.Provisional Patent Application Ser. No. 61/048,453, filed on Apr. 28,2008, and U.S. Provisional Patent Application Ser. No. 61/051,445, filedon May 8, 2008, and further, is a continuation-in-part of U.S. patentapplication Ser. No. 12/106,256, filed on Apr. 18, 2008, now issued asU.S. Pat. No. 7,923,801, which claims the benefit of priority under 35U.S.C. 119(e) to U.S. Provisional Patent Application Ser. No.60/912,581, filed on Apr. 18, 2007, U.S. Provisional Patent ApplicationSer. No. 60/958,846, filed on Jul. 9, 2007, U.S. Provisional PatentApplication Ser. No. 60/970,211, filed on Sep. 5, 2007, U.S. ProvisionalPatent Application Ser. No. 61/026,440, filed on Feb. 5, 2008, U.S.Provisional Patent Application Ser. No. 61/026,650, filed on Feb. 6,2008, U.S. Provisional Patent Application Ser. No. 61/028,481, filed onFeb. 13, 2008, and U.S. Provisional Patent Application Ser. No.61/046,379, filed on Apr. 18, 2008, all of which are incorporated hereinby reference in their entireties.

TECHNICAL FIELD

The disclosure herein relates generally to optical and electronicdevices, systems and methods that include optically sensitive material,such as nanocrystals or other optically sensitive material, and methodsof making and using the devices and systems.

BACKGROUND

Half of the sun's power reaching the earth lies in the infrared, butthis power is currently underutilized in large-area, low-costphotovoltaics. Early progress towards infrared solution-processedphotovoltaics has emerged in organic materials; however, low-bandgapconjugated polymer/fullerene-derivative bulk heterojunctions remainsensitive only to 1000 nanometers (nm). Several other efforts havefocused on sensitizing organic devices to the infrared using conjugatedpolymers/nanocrystal composites, but efficiencies are still well belowthe best organic bulk-heterojunction devices. Recent results have shownthat it is possible to make efficient photovoltaic devices comprisingonly nanocrystals films; the most efficient nanorod heterojunctiondevices show response to 800 nm.

Conjugated polymers have been widely investigated and have shownpromising efficiencies. However, they remain transparent in most of theinfrared spectral region. Because half the sun's energy lies in theinfrared, the optimal bandgap for a single-junction solar cell lies inthe infrared, well beyond the sensitivity of today's organic solarcells.

In contrast with organics and polymers, colloidal quantum dots (CQDs)offer tuning to access different spectral regions through simplevariation of their chemical synthesis. By virtue of their size-tunableoptical properties, lead salt colloidal quantum dots (CQD) can beengineered to access the visible and the short-wavelength infraredspectral regions. Recently, organic polymers sensitized using infraredlead salt nanocrystals have been investigated; however, these devicesdid not exceed monochromatic power conversion efficiencies of 0.1%.Relatively higher (e.g., 1.3%) monochromatic infrared power conversionefficiencies have been reported through the use of thiols and hightemperature processes to achieve smooth films on rough nanoporoustransparent metal oxides. The highest infrared monochromatic externalquantum efficiencies (EQE) achieved has been reported as 37% under 12 mWcm⁻² illumination at 975 nm. These PbS CQD-based devices registered aninfrared power conversion efficiency of 4.2%.

Solution-processed photovoltaics offer solar energy harvestingcharacterized by low cost, ease of processing, physical flexibility, andlarge area coverage. Conjugated polymers, inorganic nanocrystals (NCs),and hybrid materials have been widely investigated and optimized to thispurpose. Organic solar cells have already achieved 6.5% solar conversionefficiencies. However, these devices fail to harvest most of theinfrared (IR) spectral region. High efficiency muttijunction solar cellsoffer the prospect of exceeding 40% efficiency through the inclusion ofinfrared-bandgap materials. In this context, infrared single junctionsolar cells should be optimized for infrared power conversion efficiencyrather than solar power conversion efficiency. For double and triplejunction solar cells, the smallest-bandgap junction optimally lies at1320 nm and 1750 nm respectively. Attempts to extend organic solar cellefficiency into the near infrared have so far pushed the absorptiononset only to 1000 nm.

INCORPORATION BY REFERENCE

Each patent, patent application, and/or publication mentioned in thisspecification is herein incorporated by reference in its entirety to thesame extent as if each individual patent, patent application, and/orpublication was specifically and individually indicated to beincorporated by reference.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1A shows the device architecture comprising Al on a PbS nanocrystal(NC) film, under an embodiment.

FIG. 1B shows the energy band model of the device, under an embodiment.

FIG. 2 shows the current-voltage curve and photovoltaic performance(under 975 nm, 12 mW cm⁻² illumination) for a device, under anembodiment.

FIG. 3A is a comparison of the current-voltage characteristics for afirst cell configuration in the dark and under illumination fromvariations of the simulated solar illumination source, under anembodiment.

FIG. 3B is a comparison of current-voltage curves for a first cellconfiguration (under 975 nm illumination) and a second cellconfiguration (under 1550 nm illumination), under an embodiment.

FIG. 3C shows EQE spectra for a first cell configuration and a secondcell configuration, under an embodiment.

FIG. 4 shows a summary of the photovoltaic device performance parametersobtained from the various devices, under an embodiment.

FIG. 5A shows Transmission Electron Microscopy of as-synthesizedoleic-acid capped PbSe Nes (diameter approximately 5 nm), under anembodiment.

FIG. 5B shows Transmission Electron Microscopy of PbSe Nes followingoctylamine ligand exchange (the inter-nanoparticle distance wasreduced), under an embodiment.

FIG. 5C shows Transmission Electron Microscopy of networks of PbSe NCsafter benzenedithiol treatment due to the strong affinity of thethiol-end groups for the Pb atoms, under an embodiment.

FIG. 5D shows absorption of single treated (red) and double treated(blue) layers of PbSe NCs, under an embodiment.

FIG. 6 shows the performance of three differently processed PbSe CQDdevices, two of which exhibited similar external quantum efficiencies,under an embodiment.

FIG. 7 shows a table summarizing the best performances of differentlyprocessed devices recorded under 12 mW cm⁻² illumination at 975 nm,under an embodiment.

FIG. 8A shows photovoltaic device performance, specificallycurrent-voltage characteristics of a benzenedithiol treated two-layereddevice exhibiting 3.6% monochromatic power conversion efficiency at 975nm under 12 mW cm⁻² illumination, under an embodiment.

FIG. 8B shows photovoltaic device performance, specifically simulatedsolar power conversion efficiency of more than 1.1% (i.e. AM1.5 at 100mW cm²), under an embodiment.

FIG. 8C shows photovoltaic device performance, specifically spectralexternal quantum efficiency of a device reaching 37% in the infrared andabout 70% in the visible range, under an embodiment.

FIG. 8D shows a spatial band diagram showing the device model, under anembodiment.

FIG. 9A is a plot of EQE stability comparison of benzenedithiol treatedPbSe devices with previously reported amine-capped devices stored in airand in inert atmosphere, under an embodiment.

FIG. 9B is a plot of PLE stability comparison of benzenedithiol treatedPbSe devices with previously reported amine-capped devices stored in airand in inert atmosphere, under an embodiment.

FIG. 10A shows current-voltage characteristics of BDT treated PbSe CQDdevices with bottom ITO contact and with Au top contact, under anembodiment.

FIG. 10B shows current-voltage characteristics of BDT treated PbSe CQDdevices with bottom ITO contact and with Ag top contact, under anembodiment.

FIG. 10C shows current-voltage characteristics of BDT treated PbSe CQDdevices with bottom ITO contact and with Al top contact, under anembodiment.

FIG. 11A shows an input pulse of linearly increasing voltage (A=75000VIs) for typical CELIV transients.

FIG. 11B shows the output current density transient for typical CELIVtransients. The hole mobility is determined from t_(max).

FIG. 12 summarizes the contribution of the depletion and quasi-neutralregions to the EQE under 12 mW cm⁻² intensity at 975 nm, under anembodiment.

FIG. 13 is a schematic diagram of the analytical model used indetermining where electron-hole pairs were generated, under anembodiment.

FIG. 14 shows a representative ToF transient plot.

FIG. 15 is a plot of carrier recombination lifetime (blue, left axis)and external quantum efficiency (red, right axis) versus illuminationintensity at 975 nm, under an embodiment.

FIG. 16 shows a typical OCVD transient.

FIG. 17 shows transfer characteristics of PbSe NC thin film field-effecttransistors, under an embodiment.

FIG. 18 summarizes the calculated charge transport characteristics,under an embodiment.

FIG. 19A shows TEM images of chalcopyrite (CuGaSe₂) nanoparticiessynthesized along with their corresponding SAED pattern, under anembodiment.

FIG. 19B shows TEM images of chalcopyrite (CuInSe₂) nanoparticiessynthesized along with their corresponding SAED pattern, under anembodiment.

FIG. 19C shows TEM images of chalcopyrite (CIGS) nanoparticiessynthesized along with their corresponding SAED pattern, under anembodiment.

FIG. 20A shows powder XRD patterns of CuGaSe₂, CuInSe₂ and CIGSnanoparticies, under an embodiment. The vertical lines below indicatethe corresponding reflection peaks for bulk CuIn₀ ₅Ga₀ ₅Se₂, (JCPDS40-1488), CuGaSe₂, (JCPDS 79-1809) and CuInSe₂, (JCPDS 40-1487).

FIG. 20B shows an ensemble of UV-vis-NIR absorption spectrum of CuGaSe₂,CuInSe₂ and CIGS nanoparticies in toluene, under an embodiment.

FIG. 21A through 21C show plots of size distribution of as-synthesizednanoparticies (data based on manual counts of 80 nanoparticies from TEMimages), under an embodiment.

FIG. 22A shows TEM images of CuGaSe₂ synthesized by cooking Cu(Ac),Ga(acac)₃ and Se powder in oleylamine at 250 C, under an embodiment.

FIGS. 22B and 22C show TEM images of CuInSe₂ synthesized by cookingCu(Ac), In(Ac)₃ and Se powder in oleylamine at 250 C, under anembodiment.

FIG. 23A shows TEM images and corresponding SAED of CuGaSe₂ hexagonalmicroplates obtained in oleylamine and oleic acid mixture, under anembodiment.

FIG. 23B shows TEM images and corresponding SAED of CuInSe₂ hexagonalmicroplates obtained in oleylamine and oleic acid mixture, under anembodiment.

FIG. 24A shows TEM images of CuInSe₂ nanoparticies synthesized fromCu(acac)₂ and In(Ac)₃ precursors at 250° C. (scale bars are 50 nm),under an embodiment.

FIG. 24B shows TEM images of CuInSe₂ nanoparticies synthesized fromCu(Ac) and In(Ac)₃ precursors at 250° C. (scale bars are 50 nm), underan embodiment.

FIG. 24C shows TEM images of CuInSe₂ nanoparticies synthesized fromCu(acac)₂ and In(acac)₂ precursors at 250° C. (scale bars are 50 nm),under an embodiment.

FIG. 24D shows TEM images of CuInSe₂ nanoparticies synthesized fromCu(Ac) and In(acac)₃ precursors at 250° C. (scale bars are 50 nm), underan embodiment.

FIG. 25A shows TEM images of CIGS nanoparticies synthesized by injectionCu(acac)₂, In(acac)₃ and Ga(acac)₃ oleylamine solution intoSe/oleylamine at an injection temperature of 270° C., under anembodiment.

FIG. 25B shows TEM images of CIGS nanoparticies synthesized by injectionCu(acac)₂, In(acac)₃ and Ga(acac)₃ oleylamine solution intoSe/oleylamine at an injection temperature of 220° C., under anembodiment.

FIG. 26A shows TEM images of CIGS nanoparticles synthesized with aprecursor ratio CIGE3228-0.15 mmol Cu(acac)₂, 0.1 mmol Ga(acac)₃ and 0.1mmol In(acac)₃ to 0.4× mmol Se, under an embodiment.

FIG. 26B shows TEM images of CIGS nanoparticles synthesized with aprecursor ratio CIGE4138-0.20 mmol Cu(acac)₂, 0.15 mmol Ga(acac)₃ and0.05 mmol In(acac)₃ to 0.40 mmol Se, under an embodiment.

FIG. 26C shows XRD patterns of CIGS nanoparticies synthesized withprecursor ratios CIGE3228-0.15 mmol Cu(acac)₂, 0.1 mmol Ga(acac)₃ and0.1 mmol In(acac)₃ to 0.4 mmol Se, and CIGE4138-0.20 mmol Cu(acac)₂,0.15 mmol Ga(acac)₃ and 0.05 mmol In(acac)₃ to 0.40 mmol Se, under anembodiment.

FIG. 27 shows XRD patterns of CIGS nanoparticles arrested for differentreaction duration, under an embodiment.

FIG. 28A shows representative TEM images and SAED pattern of CuInS₂nanoparticies produced in oleylamine using sulfur powder instead ofselenium powder, under an embodiment.

FIG. 28B shows representative TEM images and SAED pattern of CuGaS₂nanoparticles produced in oleylamine using sulfur powder instead ofselenium powder, under an embodiment.

FIG. 29 shows composition of CIGS nanoparticles of an embodimentcalculated from Inductively Coupled Plasma Atomic Emission Spectrometry(ICP).

DETAILED DESCRIPTION

Large-area, low-cost light sensors and energy harvesters based onsolution-processed semiconductors are of wide interest. In both lightsensors and energy harvesters, devices have high longevity, includingunder intense illumination, and at high temperatures, are desired.Additional properties desired for light sensors include but are notlimited to the following: high photocurrent including large externalquantum efficiency (electrons of primary photocurrent per second perphoton incident per second) and, in embodiments, large gain (electronsof total photocurrent per second per photon incident per second); lowdark currents; a high ratio of photocurrent to dark current for a givenlevel of illumination; rapid temporal response compatible with imaging.Additional properties desired for energy harvesters include but are notlimited to the following: efficient conversion of optical power of aparticular wavelength into electrical power (a high monochromatic powerconversion efficiency); efficient conversion of a band of wavelengthsinto electrical power (in the case of the totality of the sun'sspectrum, a high solar AM1.5 power conversion efficiency); relativelylarge external quantum efficiency; relatively large open-circuitvoltage; relatively large fill-factor.

The embodiments described herein include the realization of materialshaving the aforementioned properties. Embodiments described hereininclude a composite material comprising semiconductor nanocrystal andorganic molecules. The organic molecules passivate the surfaces of thesemiconductor nanocrystals, and facilitate the transfer of chargebetween the semiconductor nanocrystals. Enhancements in the mobility ofat least one type of charge carrier (e.g., electrons, holes, bothelectrons and holes) are achieved in an embodiment via delocalization ofat least one type of carrier across at least a portion of the organicmolecule employed in passivation. In embodiments, a p-type semiconductormaterial comprises semiconductor nanocrystals where the mobility ofholes is greater than or equal to the mobility of electrons. Inembodiments, at least one benzene ring forms at least a portion of theorganic molecules and provides delocalization of at least one type ofcharge carrier, such as electrons, thereby facilitating the transport ofthat type of charge carrier.

The embodiments described herein further include the realization ofdevices that meet the aforementioned properties. In embodiments, asemiconductor material is electrically addressed using a first and asecond electrode. In embodiments, the semiconductor material is a p-typesemiconductor comprising semiconductor nanocrystals, wherein themobility of electrons in the semiconductor material is greater than orequal to the mobility of holes. In embodiments, a device, comprising afirst and a second electrode, as well as a semiconductor materialcomprising semiconductor nanocrystals, provides for the sensitivedetection of light, including a combination of high external quantumefficiency, low dark current, and video-frame-rate-compatible temporalresponse. In embodiments, the device may further provide gain, whereinmore than one electron of current flows per second for every photon persecond of illuminating light. In embodiments, a device, comprising afirst and a second electrode, as well as a semiconductor materialcomprising semiconductor nanocrystals, provide efficient conversion ofoptical power into electrical power.

Materials from which the semiconductor nanocrystals of an embodiment aremade may include one or more of the following, but the embodiment is notso limited: PbS; PbSe; PbTe; CdS; CdSe; CdTe; SnS; SnSe; SnTe; Si; GaAs;Bi2S3; Bi2Se3; CuInS2; CuInSe2; Cu(InGa)Se2 (CIGS); CuGaSe2.

Materials incorporated into the organic component of the film of anembodiment may include one or more of the following, but the embodimentis not so limited: Benzenedithiol; Dibenzenedithiol; Mercaptopropionicacid; Mercaptobenzoic acid; Pyridine; Pyrimidine; Pyrazine; Pyridazine;Dicarboxybenzene; Benzenediamine; Dibenzenediamine.

In the following description, numerous specific details are introducedto provide a thorough understanding of, and enabling description for,embodiments of the invention. One skilled in the relevant art, however,will recognize that these embodiments can be practiced without one ormore of the specific details, or with other components, systems, etc. Inother instances, well-known structures or operations are not shown, orare not described in detail, to avoid obscuring aspects of the disclosedembodiments.

Schottky-Quantum Dot Photovoltaics for Infrared Power Conversion

The embodiments described herein provide planar, stackable PbSnanocrystal quantum dot photovoltaic devices with infrared powerconversion efficiencies up to 4.2%. This represents a three-foldimprovement over the previous efficiencies obtained in the morecomplicated, stacking-incompatible nanoporous architecture. The planarSchottky photovoltaic devices described herein were prepared fromsolution-processed PbS nanocrystal quantum dot films with aluminum andindium tin oxide contacts. These devices exhibited up to 4.2% infraredpower conversion efficiency, which is an approximate three-foldimprovement over previous results, and solar power conversion efficiencyreached 1.8%. The architecture of the devices described herein allowsfor direct implementation in multi junction photovoltaic deviceconfigurations.

Current-voltage characteristics of the devices of an embodiment weremeasured in air with an Agilent 4155C Semiconductor Parameter Analyzer.Diode lasers operating at 975 and 1550 nm were used for monochromaticillumination, while solar illumination at AM1.5 conditions was simulatedwith an Oriel solar simulator (Xe lamp with filters). The sourceintensities were measured with a Melles-Griot thermopile power meter(calibration uniform within ±5% from 300 to 2000 nm), through a circular3.1 mm² aperture at the position of the sample. The external quantumefficiency (EQE) spectra were measured with a Keithley 6430Source-Measure Unit in current sensing mode, while illumination wasprovided by a white light source dispersed by a Jobin-Yvon Triax 320monochromator.

Nanocrystal films were spin-coated in an inert atmosphere onto indiumtin oxide (ITO)-coated glass substrates from a 150 mg mL⁻¹ octanesolution to produce films between 100 and 300 nm thick. The Schottkycontact was fanned using a stack of 0.7 nm LiF/140 nm Al/190 nm Agdeposited by thermal evaporation through a shadow mask; all devices hada top contact area of 3.1 mm². Photocurrents were seen to scale withdevice area up to 7.1 mm², and negligible photocurrents were observedwhen the contacts were illuminated from the metallized side.

FIG. 1A shows the device architecture comprising Aluminum (Al) on a PbSnanocrystal (NC) film (the inset shows a SEM of the nanocrystal film),under an embodiment. The device comprises a first electrode that is thetransparent conducting ITO contact. The device further comprises asecond electrode that includes the Al. The device includes asemiconductor layer positioned between the first and second electrodes.The semiconductor layer of an embodiment includes colloidal quantum dots(CQDs) passivated with organic ligands, such as PbS CQDs passivated withbutylamine, or passivated with benzenedithiol. Spin-coating thenanocrystals from octane solutions led to smooth, densely-packed arrays,as shown by the scanning electron micrograph (SEM) of the inset (scalebar is 20 nm).

FIG. 1B shows the energy band model of the device, under an embodiment.A Schottky barrier was formed at the junction between thermallydeposited Al and the ptype PbS colloidal nanocrystal film, and was theelectron-extracting contact. Photogenerated holes were extracted throughthe transparent conducting ITO contact. The energy band modelillustrates the presence of bending in the conduction band (E_(c)),valence hand (E_(v)), and vacuum energy (E_(vac)) near theAl/nanocrystal interface, under an embodiment. Photogenerated electron(e⁻) and hole (h⁺) transport is governed by the presence of a built-inelectric field within the depletion layer (of width W) of thenanocrystal layer (of thickness d). The Fermi level (E_(F)) is drawn toshow the p-type conduction characteristics. The bandgap (E_(g)) of thesenanocrystals is approximately 0.75 eV, defined by the first maximum inthe absorption spectrum.

Through optimization of the material processing steps, the photovoltaicperformance of this structure was improved. The PbS nanocrystals weresynthesized using an organometallic route and ligand-exchanged ton-butylamine as described previously. FIG. 2 shows the current-voltagecurve and photovoltaic performance (under 975 nm, 12 mW cm⁻²illumination) for a device processed using the optimized passivation(curve 2) procedure compared to the baseline (curve 1), showing anincrease in V_(oc), under an embodiment. Nanocrystals having undergoneboth the optimized passivation and ligand exchange procedures (curve 3)yielded devices with enhanced V_(oc) and EQE compared to the baseline;this is emphasized by the enclosure that represents the maximum power(P_(m)) load conditions. Fill factor is represented as “FF” and powerconversion efficiency is represented as “PCE”.

Device open-circuit voltage (V_(oc)) was increased by a factor of twowithout lowering EQE by increasing the cooling rate of the nanocrystalswith an ice bath immediately following the growth stage. Thisperformance enhancement was attributed to improved nanocrystalpassivation and hence a lower density of trap states in thesemiconductor near the metal interface, which are known to affectSchottky devices. To further improve the quality of the nanocrystal/Aljunction, a thin LiF layer was evaporated atop the nanocrystal surfaceprior to Al deposition.

As synthesized, the nanocrystals were passivated with ˜2.5 nm longoleate ligands. These long ligands prevent close nanocrystal packing andtherefore impede charge transport in films. Many of the long oleateligands were removed via repeated precipitations using methanol; asingle precipitation was used for the baseline device. To fill any emptycoordination sites and further displace oleate ligands from the surface,the process of an embodiment included a solution-phase ligand exchangeto the ˜0.6 nm long n-butylamine ligand. Nuclear magnetic resonancespectroscopy confirmed an increase in the extent of ligand exchangeincreased with the number of methanol precipitations; threemethanol-induced precipitations proved to be an optimal compromisebetween charge transport efficiency and colloidal stability. Thissolution-phase approach, in contrast with solid state ligand exchanges,enabled the spin-coating of smooth, crack-free films necessary forhigh-yield, large-area devices.

When coupled with the optimized passivation described above the exchangeenhancement resulted in a four-fold increase in power conversionefficiency. The photovoltaic performance of a representative device thathas undergone this series of procedures is shown in FIG. 2 describedabove.

These optimizations allowed the fabrication of a variety of highlyefficient solution-processed photovoltaic cells. A first cellconfiguration of an embodiment used 230 nm thick layers of PbSnanocrystals with a first excitonic transition at 1650 nm, which isclose to the optimal infrared band gap (1750 nm) in a triple junctioncell. A second cell configuration of an embodiment provided maximumpower conversion at 1550 nm, using smaller nanocrystals (with enhancedabsorption at this wavelength) in slightly thicker films (250 nm).

FIG. 3A is a comparison of the current-voltage characteristics for afirst cell configuration in the dark and under illumination fromvariations of the simulated solar illumination source, under anembodiment. FIG. 3B is a comparison of current-voltage curves for afirst cell configuration (under 975 nm illumination) and a second cellconfiguration (under 1550 nm illumination), under an embodiment. FIG. 3Cshows EQE spectra for a first cell configuration and a second cellconfiguration, under an embodiment. The inset shows the absorptionspectra of the ligand-exchanged nanocrystals in solution; excitonicpeaks were also visible in the film absorption spectra. au, arbitraryunits.

As stand-alone cells, cells of the first cell configuration yielded amaximum power conversion efficiency of 1.8% under simulated solarillumination at 100 m W cm⁻². These devices showed up to 4.2% efficiencyunder monochromatic 975 nm illumination at 12 mW cm⁻². To demonstratethe capabilities of infrared power conversion and to simulate theconditions within multijunction cells, the simulated solar source wasfiltered using amorphous Si (a-Si) and GaAs. The a-Si sample transmittedwavelengths longer than 640 nm to produce an infrared intensity of 44 mWcm⁻² incident on the device. The device converted the resultanttransmitted broadband power with 1.8% efficiency. The GaAs filtertransmitted 25 mW cm⁻² at wavelengths longer than 910 nm onto thedevice; the measured near infrared power conversion was then 1.3%. Thissuggests that these photovoltaic cells could be directly stacked withother solution-processed cells to achieve efficiencies beyond the 6%mark.

Monochromatic power conversion efficiencies of up to 2.1% were obtainedwith cells of the second cell configuration using a 1550 nm illuminationsource. This wavelength is appealing for wireless power distributionapplications. The power conversion remained close to 2% even above 100mW cm⁻² illumination, which is promising for the high intensitiesrequired in thermophotovoltaic systems.

FIG. 4 shows a summary of the photovoltaic device performance parametersobtained from the various devices, under an embodiment. While the V_(oc)values are not large in absolute terms, they are high as a fract ion ofthe material bandgap. Device performance was maintained for up to 5hours in air but degraded completely over 24 hours. Nanocrystal filmswere stable in an inert environment for long durations (seven days),which indicates that encapsulation could improve long-term stability.Short-circuit current density is represented as “J_(sc)”, and powerconversion efficiency is represented as “PCE”.

The EQE spectrum of devices was measured by illuminating the deviceswith monochromatic light, measuring the current under short-circuitconditions, and scaling them with respect to the previously indicatedEQEs measured at 975 nm and 1550 nm (for the first and second cellconfigurations, respectively). The EQE spectra for devices of bothdesigns are shown in FIG. 3C. The features of the colloidal nanocrystalabsorption spectrum (inset) are manifest in the EQE spectra wherein awell-defined first excitonic transition is retained even indensely-packed, conductive films. The EQE reached values near 60% forvisible wavelengths. Film absorption data was also acquired by measuringthe fraction of light reflected through the substrate and correcting forITO and Al absorption. The film absorptions at 975 nm and 1550 nm were41% and 14%, respectively, which indicates that the internal quantumefficiency exceeded 90% for the best devices.

Efficient, Stable Infrared Photovoltaics Based on Solution-CastColloidal Quantum Dots

The embodiments described below provide efficient solar cells (e.g.,3.6% power conversion efficiency) operating in the infrared to evincemulti-week stability. Since ligands previously employed inrecord-setting solar cells are labile and reactive with adjacent metalcontacts, militating against long-lived device efficiency, astrongly-binding end functional group was selected to passivate thenanoparticle surfaces robustly in the solid state of an embodiment,while avoiding reactivity with the adjacent metal contact. A furtherincrease in proximity among the nanoparticles could be achieved, andcould result in improved electron and hole transport, withoutsacrificing the highly desired quantum size-effect tuning offered by theuse of colloidal quantum dots, so the devices described herein comprisea short bidentate linker having a conjugated, instead of an entirelyinsulating, moiety lying between the end groups. Because a large achange in film volume resulting from the exchange of longeroleic-acid-capped ligands to short crosslinkers would lead to poor filmmorphology and electrical shorts, solution-exchange to a shorter linkerwas used in an embodiment prior to film formation and film crosslinkingThese considerations, taken together, led to use of a solution-phaseexchange to an intermediate ligand, octylamine, followed bysolution-casting of films, and finished with a treatment using thebidentate linker, benzenedithiol.

In contrast, the PbSe CQD based photovoltaic device of an embodimentachieves 3.6% infrared power conversion efficiency (PCE). This appearsto be the first PbSe colloidal quantum photovoltaic to exceed onepercent infrared PCE. It also represents the first solution-processedinfrared photovoltaic device to be stable over weeks without requiringfresh deposition of its top electrical contact. Thus, the devicesdescribed herein appear to be the first to manifest, and indeed exploit,a very surprising feature—the diffusion of charge carriers, within acolloidal quantum dot solid, over hundreds of nanometers.

The PbSe NCs of an embodiment are synthesized using a modified versionof the organometallic route reported for PbS NCs. FIG. 5A showsTransmission Electron Microscopy of as-synthesized oleic-acid cappedPbSe NCs (diameter approximately 5 nm), under an embodiment. Theas-synthesized NCs were capped with approximately 2 nm oleate ligands,previously reported to impede efficient charge transport in films. Thebenzenedithiol crosslinking was preceded with a solution-phase ligandexchange. As a result, the oleate ligands were replaced with shorteroctylamine ligands (˜1 nm). FIG. 5B shows Transmission ElectronMicroscopy of PbSe NCs following octylamine ligand exchange (theinter-nanoparticle distance was reduced), under an embodiment. FIG. 5Cshows Transmission Electron Microscopy of networks of PbSe NCs afterbenzenedithiol treatment due to the strong affinity of the thiol-endgroups for the Pb atoms, under an embodiment.

Regarding the benzenedithiol crosslinking, the crosslinking processperformed on the CQD films of an embodiment was optimized by varying thetreatment durations through a 1,4-Benzenedithiol treatment optimization.After exposing the NC layers to BDT, films were blow dried with nitrogenand subjected to 30 min vacuum drying. The crosslinking treatment wasfurther optimized by varying its duration time for each of the twolayers. Typically, a 20 min treatment time was used for both layers;however, devices where the first layer was treated for shorter durations(˜10 min) while the second one for longer (˜30 min) recorded the highestexternal quantum efficiencies, reaching 46% and monochromatic powerconversion efficiencies of more than 3.6% under 12 mW cm⁻² at 975 nm.The reduced treatment time for the first layer enabled realization ofsmoother and defect-free films. In addition, the fill factor and theopen circuit voltage both degraded when the first layer was subjected tomore than 15 min treatment time. The later fact correlates with thereduction of the shunt resistance of the device, i.e. pinholes due tothe chemical processing are significant.

FIG. 6 shows the performance of three differently processed PbSe CQDdevices, two of which exhibited similar external quantum efficiencies,under an embodiment. The best passivated device 604 exhibits an IR PCEof 3.65% with 40% fill factor. The fill factor decreases to 34% when thetreatment duration of the first layer 606 is increased. The highest fillfactor (45%) is registered for a reduced treatment time for both layers602. The device with 20 min treatment duration for both layers 606 wasdeeply affected by processing defects because it exhibited the lowestopen circuit voltage (V_(oc)) and fill factor. Optimization of thetreatment duration was therefore essential to maximizing the transportproperties and thus the performance while minimizing pinholes. FIG. 6also shows the best performance of an under-treated device 602 which hada better fill factor resulting from minimal defects but was not asefficient as the best passivated devices 604. FIG. 7 summarizes the bestperformances of differently processed devices recorded under 12 mW cm⁻²illumination at 975 nm, under an embodiment.

The embodiments herein used PbSe NCs having an absorption peak rangingbetween 1200 and 1300 nm. FIG. 5D shows absorption of single treated 502and double treated 504 layers of PbSe NCs, under an embodiment. Thin NCfilms (˜110 nm) were spin-coated on ITO substrates and the samples wereimmersed in a dilute solution of benzenedithiol in acetonitrile (3.5 mM)for a duration ranging from 10 to 30 minutes. This rendered the layerinsoluble in the nonpolar solvents that were used for spin-coating theNCs. A second thin layer was deposited on top to ensure the formation ofa smooth, densely packed film. The second layer was also subjected to alinking treatment. The total thickness of the NCs active layer rangedbetween 210 and 250 nm.

FIG. 8A shows photovoltaic device performance, specificallycurrent-voltage characteristics of a benzenedithiol treated two-layereddevice exhibiting 3.6% monochromatic power conversion efficiency at 975nm under 12 mW cm⁻² illumination, under an embodiment.

FIG. 8B shows photovoltaic device performance, specifically simulatedsolar power conversion efficiency of more than 1.1% (i.e. AM1.5 at 100mW cm⁻²), under an embodiment. From total absorbance measurements at 975nm, the IQE was found to approach 90%.

FIG. 8C shows photovoltaic device performance, specifically spectralexternal quantum efficiency of a device reaching 37% in the infrared andabout 70% in the visible range, under an embodiment. The spectrallyresolved EQE is presented between 400 to 1600 nm. The EQE followsclosely the features of the absorption spectrum shown in FIG. 1D; awell-defined first excitonic peak is observable at 1250 nm. In thevisible wavelengths, a peak: EQE of 70% is recorded. From measurementsof total film absorbance, the internal quantum efficiency is estimatedat 975 nm to approach 90% in the best devices, which implies highlyefficient charge separation and extraction. With respect toreproducibility of the results reported herein, throughout the course ofthis study, approximately 40 devices were made that exhibited infraredmonochromatic power conversion efficiencies in excess of 3%.

FIG. 8D shows a spatial band diagram showing the device model, under anembodiment. A Schottky barrier is formed at the Mg/p-type semiconductingNCs interface. The majority of the photogenerated carriers diffusethrough the quasi-neutral region (L_(QN) ⁻ 145 nm and are separated inthe depletion region (W˜65 nm). A fraction of the carriers is lost torecombination.

FIG. 9A is a plot of EQE stability comparison of benzenedithiol treatedPbSe devices with previously reported amine-capped devices stored in airand in inert atmosphere, under an embodiment. FIG. 9B is a plot of PCEstability comparison of benzenedithiol treated PbSe devices withpreviously reported amine-capped devices stored in air and in inertatmosphere, under an embodiment. This data compares the stability of thedevices fabricated as described herein with previously reportedhigh-efficiency devices (e.g., devices fabricated by spin-coatingbutylamine-capped PbS NCs on ITO substrates and evaporating Al contactson top).

The benzenedithiol-crosslinked PbSe NC devices of an embodiment retainedtheir high EQE values for over ten days when stored in a nitrogen filledglovebox (solid red line), and their PCE maintains 90% of its initialvalue for more than 2 days. The mine-capped devices (dashed red line)severely deteriorated within the first 24 h (e.g., lost half of theirEQE and more than 75% of their power conversion efficiency).

In air, the benzenedithiol treated devices (solid blue lines) registeredgreater stability than the amine-capped devices (dashed blue lines)(e.g., the dithiol-capped PbSe NC based devices retained their high EQEand ˜80% of their PCE over 48 hours, whereas the amine-capped deviceslost all performance within the same period of time). Note that all thetesting was done in air, and all the devices in this study were exposedto a minimum of 15 hours of oxygen and moisture. Several high efficiencybenzenedithiol treated devices retained over 75% of their EQE for morethan two weeks. The reported EQE and PCE values were taken under 12 mWcm⁻² at 975 nm.

The physical mechanisms responsible for the performance of the devicesof an embodiment was investigate in order to understand further theprocess steps that were most critical to the performance improvementsnoted in the devices. The investigation, described in detail below,began by determining which of the metal-semiconductor junctions isresponsible for providing rectification and charge separation. Estimateswere made from capacitance-voltage measurements of the spatial extent ofthe resultant depletion region and, from this, a spatial band diagramwas proposed. The absorbance of photons was quantified in each of thekey regions (the depletion and quasi-neutral portions) of the colloidalquantum dot active region. Combining this with knowledge of the depth ofthe depletion region, the thickness of the device, and the measured EQE,estimates were made of the minority carrier diffusion length required toexplain the relatively high overall efficiency. An investigation wasthen carried out as to whether the electronic properties of thematerials of an embodiment are indeed capable of supporting such a longminority carrier diffusion length. Electron mobility was measured usingthe time-of-flight method and the carrier lifetime was estimated usingthe open-circuit-voltage decay method. In this way an electron diffusionlength was estimated in the range 200-300 nm, and this range was foundto be sufficient to account for the relatively high observedefficiencies. It was also concluded that further improvement inperformance could be achieved by increases in electron mobility leadingto even more efficient diffusion from the quasi-neutral region to theedge of the depletion region.

The location of the rectifying metal-semiconductor junction, and thusthe depletion region, in the devices of an embodiment was studied byvarying the choice of metal both as bottom and top contact. First, whilemaintaining ITO as the bottom contact, Al (˜4.2 eV), Ag (˜4.3 eV), or Mg(˜3.6 eV) was deposited atop the NC films. In all cases similar I-Vcharacteristics were obtained, though the open circuit voltages (V_(OC)_(—) _(Mg)>V_(OC) _(—) _(Al)>V_(OC) _(—) _(Ag)) were smaller when Ag andAl were used. On the other hand, when Au was employed as the topcontact, the device exhibits a linear I-V (see FIG. 10A). From this itwas concluded that there exists a depletion region at the Mg—NC junctionin the devices. FIG. 10A shows current-voltage characteristics of BDTtreated PbSe CQO devices with bottom ITO contact and with Au topcontact, under an embodiment. FIG. 10B shows current-voltagecharacteristics of BDT treated PbSe CQD devices with bottom ITO contactand with Ag top contact, under an embodiment. FIG. 10C showscurrent-voltage characteristics of BDT treated PbSe CQD devices withbottom ITO contact and with Al top contact, under an embodiment.

The bottom contact was also varied, replacing ITO (˜4.8 eV) with Pt(˜6.3 eV) or Au (˜5 eV). All devices continued to provide goodrectification when a Mg top contact was employed; and thecurrent-voltage characteristics did not appreciably change. From this itwas concluded that the bottom contact (transparent ITO the devices)serves mainly for ohmic hole collection in the solar cells. The spatialband diagram shown in FIG. 8D is proposed based on these conclusions.

The depth of the depletion region was then determined. The capacitancewas measured at zero bias (under short-circuit conditions). The staticrelative permittivity via the charge extraction was found by linearlyincreasing voltage (CELIV) method to be 15±1.

The majority carrier mobility in the photovoltaic devices were measuredvia CELIV, which involves applying a pulse of linearly increasingvoltage (U(t)=At, where A is the pulse slope) and measuring the currenttransient across the sample with at least one blocking contact. Usingthe CELIV method, the extraction of the equilibrium carriers wasinvestigated. The time to reach the extraction current maximum t_(max)was used to estimate the hole mobility. When Δj=j_(max)−j(0)≦j(0), wherej_(max) is the maximum value of the current and j(0) is the displacementcurrent as indicated in figure S6, the hole mobility can be calculatedfrom

$t_{\max} = {d\sqrt{\frac{2}{\mu\; a^{\prime}}}}$

where d is the film thickness (˜230 nm), and μ is the majority carriermobility. The mobility was found to be 2.4×10⁻³ cm² V⁻¹s⁻¹ in the PbSeNC based solar cells of an embodiment.

The displacement current step j(0) can be used to evaluate the staticrelative permittivity ∈_(r),j(0)=∈_(r),∈₀ A/d

where ∈₀ is permittivity of free space. The static relative permittivitywas found to be 15±1. A typical CELIV transient is represented in figureS6. FIG. 11A shows an input pulse of linearly increasing voltage(A=75000 V/s) for typical CELIV transients. FIG. 11B shows the outputcurrent density transient for typical CELIV transients. Hole mobility isdetermined from t_(max).

This allowed for estimation of the depletion width to be 65±5 nm. Thecapacitance was measured at zero bias under short circuit conditions inorder to determine the device depletion width (W). The capacitance valueper unit area (C_(i)) was determined to be 2×10⁻⁷ F cm⁻². The depletionwidth was determined using the following equationW=∈ _(r)∈₀ /C _(i)where ∈₀ is permittivity of free space, and ∈_(r) is static relativepermittivity.

Having elucidated a simple spatial band diagram, efforts turned todetermining where electron-hole pairs were generated, and in whichquantities, within the two regions of interest—the depletion region,with its charge-separating field, and the quasi-neutral region, in whichminority carrier diffusion would serve as the dominant transportmechanism.

The analysis first looked at whether absorption within the depletionregion could, on its own, account for the high observed EQEs. Fromknowledge of the absorption per unit length of 975 nm light, thefraction of incident power absorbed in each region was determined, andis shown in FIG. 12. This led to the conclusion that less than half ofthe observed short-circuit current is attributable to electron-hole pairgeneration from absorption within the depletion region.

To account for the EQE observed, it is estimated that two thirds of theelectron-hole pairs photogenerated within the quasi-neutral portion ofthe device have diffused to the depletion region to be efficientlyseparated therein and extracted therefrom. Such a result would bepossible if the minority carrier diffusion length for electrons in thequasi-neutral region were to exceed a few hundred nanometers.

For the total absorption measurements, the reflectivity of the substratewas measured and corrected for the ITO (˜5%) and Mg contact absorption(˜4%). In the case considered in the following analysis, the percentageof light absorbed at 975 nm was 43% and the corresponding mean externalquantum efficiency was 32%. It is assumed that the light absorbed in thephotovoltaic devices through the ITO contact was reflected back at themirror-like upper metallic contact.

Thus, assuming a uniform light absorption profile, the fraction of lightabsorbed incident from the ITO side in double pass is:A _(total)=1−e ^(−2ad)

where d is the thickness of the CQD film, and α is the absorbance. Thepercentage of photons absorbed through the depletion region of 65 nmwidth near the Mg contact in double-pass is estimated to be 13%.Accordingly, the electron-hole pairs generated in the depletion region,believed to be efficiently extracted (>90%), only contributed to lessthan half of the photocurrent. The quasi-neutral region absorbed theremaining 30% of the incident light. In order to account for the 32% EQEat 975 nm, 20% of the photogenerated carriers created in thequasi-neutral portion of the device must have diffused to the depletionregion to be separated therein. The rest was lost to carrierrecombination. Diffusion therefore must be playing a large role in thedevices of an embodiment.

FIG. 13 is a schematic diagram of the analytical model used indetermining where electron-hole pairs were generated, under anembodiment The electron-hole pairs generated within the depletion region(W) are efficiently separated. The electrons generated within L_(QN2)diffuse to the depletion region where they are separated by the built-infield while the electrons generated within L_(QN1) of the ITO contactare mostly lost to recombination. Assuming that the electron-hole pairswhich are nearest to the depletion region have a larger chance todiffuse and separate, the photons absorbed in L_(QN2) of thickness ˜95nm account for the rest of the EQE. In order to estimate the requireddiffusion length for the least mobile carrier (electrons in this case),it is assumed that only the photogenerated carriers in the quasi-neutralregion that have a transit time of 0.1τ in the depletion region(L_(QN2)˜95 nm) are largely not lost to recombination where τ is therecombination lifetime. The model is estimated to be plausible withinexperimental uncertainty if the minority carrier diffusion length is inthe 200-300 nm range.

To evaluate the feasibility of the electron and hole drift over thedepletion region depth (˜65 nm) and the electron minority carrierdiffusion over the quasi-neutral region depth (˜145 nm), the electronand hole mobilities were measured. When combined with carrier lifetimes,these enable the drift and diffusion lengths to be estimated.

Minority electron carrier mobility was studied using the time-of-flight(TOF) method and the majority carrier mobility via charge extraction bylinearly increasing voltage (CELIV). TOF experiments employed a samplewith a geometry identical to the photovoltaic device, i.e. a layer ofNCs sandwiched between the ITO and magnesium contacts, with theexception that the total NC layer in this case was thicker (>500 nm).The electron mobility was found to be 1.4×10⁻³ cm²V⁻¹s⁻¹ for abenzenedithiol treated device. CELIV experiments conducted on thephotovoltaic devices allowed us to estimate the hole mobility to be2.4×10⁻³ cm²V⁻¹s⁻¹ in the dithiol treated NC based devices. Thus, theelectron and the hole mobility in the devices of an embodiment arewithin the same order of magnitude, in contrast with recent findings inPbS colloidal quantum dots devices, where the minority electrons were ˜8times less mobile.

Time of flight (ToF) was performed on a sample with a geometry identicalto the photovoltaic device, i.e. a layer of NCs sandwiched between theITO and magnesium contacts, with the exception that the total NC layerin this case was thicker (>700 nm). The devices were held under reversebias by applying a positive potential to the Mg contact and a negativepotential to the ITO. A 10 ns pulse of 532 nm light was incident on thesample from the transparent ITO side. Under the influence of the appliedbias, the photo-generated minority carriers (electrons) drift across thesample to the Mg anode. The transient current generated by the flow ofthe photocarriers was recorded. In the case of dispersive media such asNCs, the transit time was determined by finding the intersection of twolinear regions on a log-log plot of current versus time. FIG. 14 shows arepresentative ToF transient plot. The intersection of the two linearregions 1402 and 1404 is the transit time t₁˜1 μs. The sample thicknesswas 750 nm and the voltage bias was set to 5 V. The transit time isrelated to the mobility by μ=d²/V.t_(r) where d is the film thicknessand V is the applied bias. The minority electron mobility registered twoorders of magnitude increase from almost 1×10⁻⁵ cm²V⁻¹s⁻¹ for anuntreated film to 1.4×10⁻³ cm²V⁻¹s⁻¹ for a BDT treated film.

The recombination lifetime τ was estimated at relevant solar intensitiesthrough the technique of transient open circuit voltage decay (OCVD).The device was illuminated using a digitally modulated 975 nm diodelaser at different intensities. At 12 mW cm⁻² the lifetime was found tobe on the order of 10-20 μs. FIG. 15 is a plot of carrier recombinationlifetime (dots, with reference to left axis) and external quantumefficiency (squares, with reference to right axis) versus illuminationintensity at 975 nm, under an embodiment. The decrease in the EQE (>10mW cm⁻²) corresponds to the limit where the minority carrier transittime exceeds the recombination lifetime.

In order to efficiently extract carriers, the transit time should beshorter than the characteristic time for carrier relaxation at therelevant intensities, which was measured using the technique oftransient open circuit voltage decay (OCVD). The OCVD method wasperformed on a photovoltaic device by abruptly turning off theillumination and recording the V_(oc) decay. The device was illuminatedusing a 975 nm laser at different intensities. The voltage generatedacross the device was recorded as the illumination was abruptly removed(within 3 μs). The recombination lifetime at every intensity wasdetermined by applying a linear fit to the initial V_(oc) decay.

FIG. 16 shows a typical OCVD transient. The recombination was evaluatedfrom the following relation

$\tau = {C \times \frac{{kT}\mspace{14mu} 1}{q{{\mathbb{d}V_{oc}}/{\mathbb{d}t^{\prime}}}}}$

where k is the Boltzmann constant, T is temperature, q is the elementarycharge. The recombination lifetime is evaluated from the slope 1602 ofthe voltage transient decay. From the slope 1602, the lifetime wasestimated to be 10 μs at 20 mW cm⁻². The coefficient C varies from 1 inlow injection regime to 2 in high injection regime. For this analysis,the recombination lifetime was underestimated and set C to 1. At theoperating intensity of 12 mW cm⁻², the lifetime was found to be 13 μs.

Thin film field-effect transistors (FET) were fabricated on highlyconductive silicon wafers with 100 nm of thermally grown oxide as thegate dielectric. The source and drain electrodes were separated by a 10μm gap. In order to spin-coat 50 nm thick films, the octylamine-cappedPbSe NCs was diluted to a concentration of 10 mg mL⁻¹ and thebenzenedithiol solution to 1 mM. For constant drain source voltage, arange of gate voltages was applied and recorded the current modulationthrough the NC film.

FIG. 17 shows the FET transfer characteristics of PbSe NC thin filmfield-effect transistors, under an embodiment. These FET transfercharacteristics are for NC films directly deposited from solution andafter exposure to benzenedithiol treatment. The I_(d)−V_(d)characteristics for both films are shown in the inset. The PbSe NC filmsexhibit p-type behaviour before (dots 1702) and after benzenedithioltreatment (squares 1704). The conductivity, which can be induced fromthe slope of the I_(d)−V_(d) curves, increased after treatment. In bothcases, the drain current (I_(d)) increases in magnitude as the appliedgate voltage (V_(gs)) becomes more negative. This suggests the formationof a p-channel in the FET devices indicative of the p-type behaviour ofthe PbSe NCs.

Thin films field-effect transistors (FET) were fabricated on highlydoped n-type silicon wafers with 100 nm of thermally grown oxide as thegate dielectric. The source and drain electrodes were separated by a 10um gap. The FET devices were completed by spin-coating thin layers ofPbSe NCs. The PbSe CQD films exhibited p-type behavior both before andafter treatment.

The conductance (G) of the NC films, equated to the slope of theI_(d)−V_(d) graph at zero gate bias (inset of FIG. 15), increased afterthe film treatment. The field effect devices presented suffered fromvery low I_(on)/I_(off) ratio (weak modulation) which is a consequenceof high leakage current in the off-state, inherent of the thin film FETtransistor architecture.

Carriers in the depletion region are separated via the action of thebuilt-in field resultant from the metal-semiconductor, or Schottky,junction. The drift length is given by

$\frac{{\mu\tau}\;{Vbi}}{W},$where, μ is the carrier mobility, V_(bi) is the built in potential, andW is the depletion width. Under short circuit conditions at 12 mW cm⁻²and assuming a built in voltage of 0.3 V, drift lengths of 8.5 μm wereestimated for electrons and 14.5 μm for holes. In sum, no difficulty isexpected in removing each carrier type from the 65 nm thick depletionregion.

In the quasi-neutral region, charge transport occurs mainly throughdiffusion and the carrier diffusion length may be obtained from

$\sqrt{{\mu\tau}\frac{k\;\tau}{q}}.$The calculated electron minority diffusion length is in excess of 220nm, which allows a substantial fraction of the minority carriers todiffuse out of the neutral region, and allows accounting for the highobserved EQE. FIG. 18 summarizes the calculated charge transportcharacteristics, under an embodiment.

From the transport parameters calculated, the hole diffusion length wasestimated to be 350 nm and the electron minority diffusion length wasestimated to be 220 nm (equations included in the main text). Thesevalues are evaluated in low injection mode with the minimal value of therecombination lifetime 13 μs; thus, the minimum value of the minoritydiffusion length is 220 nm. Referring to FIG. 13, the latter value isfound to be very close to the required diffusion length.

Given a diffusion length of 220 nm and assuming that after 0.1τ a largefraction of the carriers are lost to recombination, 14% of thephotogenerated carriers diffuse through the quasi-neutral region and getseparated in the depletion region. Thus, adding the depletion regioncontribution a total of 27% EQE can be accounted for, and that is withlower limit for the recombination lifetime. Thus, the model describedherein is presumed valid and the diffusion of electrons through thequasi-neutral region to the depletion region contributes significantlyto the charge extraction process in the devices of an embodiment.

Carrier extraction in these photovoltaic devices, due to the narrowdepletion region, is critically dependent on diffusion enabled by highminority carrier mobility and long lifetime. This contrasts with recentfindings in drift-dominated PbS Schottky-barrier devices.

The proposed physical picture of an embodiment is further corroboratedby investigating the dependence of EQE and recombination as a functionof illunination intensity, as seen in FIG. 13. The EQE began to diminishat intensities greater than 10 mW cm⁻². From OCVD measurements, therecombination lifetime drops below about 10 μs at such intensities. Inview of the electron mobility, the electron diffusion length begins tocontract well below the quasi-neutral region thickness under suchconditions, accounting for the onset of EQE roll-off.

With this performance and physical picture explained, the role of thebidentate linker of an embodiment is described. The first expectedimpact of linking nanoparticles in the solid state is to bring theparticles closer together (see FIGS. 5A, 5B, and 5C). FIG. 5C shows thatbenzendithiol molecules are most likely crosslinking the PbSenanoparticles. Specifically, both electron and hole mobilities increasedby more than an order of magnitude as a result of crosslinking. Evenwith the mobility-increasing treatments, the films retained theirquantum size effect, as seen in absorption spectra and external quantumefficiency spectra.

It is noted above that diffusion plays a much larger role in the devicesof an embodiment than in the typical PbS solar cells. The electronmobility of the devices of an embodiment is approximately seven timesgreater than typical PbS solar cells, a fact which accounts for adoubling of the electron minority diffusion length in the materialssystem described herein. Benzenedithiol is in fact a molecular conductorin view of its delocalization of electron molecular orbitals. Inaddition, conjugated dithiol molecules used to bridge quantum dotsystems have previously been reported not only to link nanocrystals, butalso to provide a pathway for electron transfer.

The devices described herein exhibited a photovoltaic response onlyafter being subjected to the benzenedithiol crosslinking process. It isherein proposed that the as exchanged NCs were dominated by a largedensity of unpassivated surface states. As with previously-reportedchemical processes on PbSe and CdSe NCs, benzenedithiol offerspassivation of dangling bonds. Additionally, as seen in the stabilitystudy, benzenedithiol appears to offer a longer-lived NCs/metalinterface than do amine ligands. The latter are believed to react withthe top metal contact.

Another feature of the device processing architecture described hereinis the use of two superimposed layers of colloidal quantum dot solids toincrease absorbing thickness and minimize pinholes. Preceding thesolid-state treatment with a solution-phase exchange to a somewhatshorter ligand helped to reduce volume contraction upon crosslinking ofthe film. This contributed to the realization of densely-packed,high-mobility films in situ on a substrate.

In sum, the devices of an embodiment are stable, high-efficiencyinfrared solution-processed photovoltaic devices. Further, it was shownherein that minority carrier diffusion can occur efficiently overhundreds of nanometers in such films. Strongly-passivating, short,electron-transport-assisting bidentate linkers appear to play a key rolein achieving these properties.

The chemicals used in producing an embodiment include one or more of thefollowing, but are not so limited: Lead (II) oxide powder (PbO, 99%);Oleic acid (OA, technical grade 90%); 1-Octadecene (ODE, technical grade90%); anhydrous toluene; octane; methanol; isopropanol; acetonitrile;ethyl acetate; Bis(trimethylsilyl)selenide (TMSe); 1,4 benzenedithiol(97%).

Regarding PbSe synthesis and ligand-exchange procedures of anembodiment, ODE was degassed by pre-pumping at 80° C. for 16 hours andTMSe source was pre-filtered with 0.1 and 0.02 μm Whatman syringefilters before use. The synthesis was performed in a single, three-neck,round bottom flask. The Pb precursor was prepared by pumping the mixtureof PbO and OA at 80° C. for 16 hours. The resulting transparent solutionof lead oleate precursor was stirred vigorously while being heated underAr for about 30 min. The stock solution of selenium precursor wasprepared by mixing TMSe with ODE in a glove box and the portioncorresponding to a 2:1 (Pb:Se) molar ratio was rapidly injected into thereaction flask. The injection temperature ranged between 125° C. forsmaller nanocrystals and 140° C. for the largest NCs. Upon injection,nucleation occurs instantly; thus, rapid injection is critical toachieve a narrow size distribution. After injection, the temperature ofthe reaction was dropped down and the reaction was quenched bysubjecting it to a water-ice bath for 1 min and 40 sec. A typicalsynthesis for NCs having their excitonic peak ranging from 1200 nm and1300 nm involved injecting of 7 mL of selenium stock solution (1 mmol ofTMSe) into the reaction flask containing 2 mmol PbO (0.45 g), and 63mmol of OA. PbSe NCs, particularly in their solution phase, wereobserved to be extremely sensitive to both air and moisture and as aresult all post-synthetic treatments were performed in a glove box withanhydrous reagents. The oleate-capped PbSe NCs were isolated from anyremaining starting materials and side products by precipitating thesolution with a mixture of equal volumes of methanol (5 mL) andethylacetate (5 mL). The precipitate was then re-dispersed in tolueneand re-precipitated with methanol. After the second precipitation, theNCs were vacuum-dried for 10 min and redispersed in toluene.

The solution exchange procedure was carried out inside a nitrogen filledglovebox. The as-synthesised NCs were precipitated with methanol,vacuum-dried for 10 min, and redispersed in octylamine. After threedays, the NCs were precipitated with anhydrous isopropanol, vacuum-driedfor 10 min and redispersed in octane solution to achieve a typicalconcentration of 80 mg mL⁻¹.

Regarding device fabrication, testing, and characterization, theoctylamine-exchanged PbSe NCs were spin-coated on ITO-coated glasssubstrate inside the glovebox. The devices had a typical thickness of210 to 250 nm as measured with a surface profiler (Veeco Dektak3). Thebenzenedithiol treatment was done in a fumehood in air. Approximately100 nm Mg/190 nm Ag were deposited by thermal evaporation through ashadow mask, leading to a contact area of 3.1 mm². The devices werestored in 10a nitrogen filled glovebox for 24 h before initial testing.All device characterizations were carried out in dark shieldedenclosures in air.

All current-voltage measurements including FET characterizations weretaken with an Agilent 4155C semiconductor parameter analyzer. For the IRcharacterizations, the devices were illuminated through the ITO-coatedglass using a continuous-wave diode laser operating at 975 nm. An Orielsolar simulator operating at 100 mW cm⁻² was used to simulate the solarspectrum under AM1.5 conditions. The illumination intensity was measuredwith a Melles-Griot broadband power meter.

For the TOF measurements, thick samples (>600 nm) were excited using aYttrium-Aluminum-Gamet (YAG) laser operating at 532 nm with 10 ns pulsesat a 10 Hz repetition rate. The light was incident on the sample fromthe transparent ITO side. The devices were biased using a Keithley 2400Source Meter, and a digital oscilloscope was used to measure the currenttransient output across a 50Ω load. The CELIV measurements were carriedout using an Agilent 33120A function generator which provided thelinearly increasing voltage signals and the current output was measuredacross a 50Ω load with a Tektronix IDS 220 digital oscilloscope.

The OCVD curves were recorded using a digital oscilloscope with a 1 MS)input impedance. The illumination source (975 nm diode laser) wasmodulated using a Stanford Research Systems DG535 digital pulsegenerator. An Agilent 4284A LCR meter was used to measure thecapacitance at zero bias in order to determine the device depletionwidth.

For the external quantum efficiency spectrum measurements, the incidentlight was chopped at 100 Hz and the short-circuit current was measuredwith a Stanford Research SR830 lock-in amplifier. Illumination wasprovided by a white light source dispersed by a Jobin-Yvon Triax 320monochromator. The light intensity was kept constant for allwavelengths. The measured spectrum was then scaled to match the value ofthe monochromatic EQE obtained at 975 nm.

The total film absorbance was obtained by measuring the reflectivity ofthe substrate and correcting for the ITO and Mg contact absorption in anintegrating sphere. A Cary 500 UV-Vis-IR Scan photospectrometer in thereflective mode was used to measure the reflectivity spectra. TEM imageswere taken using a Hitachi HD-2000.

Synthesis of CIGS Nanoparticles: Achieving Monodispersity,Crystallinity, and Phase-Purity

Colloidal nanoparticles have been shown to enable the realization oflow-cost, large-area, physically flexible photodetectors, photovoltaics,optical modulators, and optical sources. Devices realized to date havetypically relied on semiconductor nanoparticles containing either Cd— orPb—. The use of these metal cations leads to a robust and facilesynthesis; however, the use of these metals raises concerns regardingcompatibility with the natural environment, especially if large-areadeployment, such as is desired in the case of solar cells, were sought.

As an alternative, the synthesis of CuInS₂ CuInSe₂ and Cu(InGa)Se₂(CIGS) nanoparticles has recently been pursued. Polycrystalline CIGSfilms have long been advanced as a solar energy conversion material,providing high efficiencies of 19.2% in a laboratory device andexcellent radiation hardness.

The realization of third-generation solar cells having greater than 30%AM1.5 power conversion efficiencies can be achieved by stackingdifferent-bandgap semiconductors on one another to form tandem (e.g.,two-junction) or multijunction cells. A first cell of the stack absorbshigher-energy photons only, and provides a large opencircuit voltage;the next cells of the stack absorb the lower-energy photons, and,series-connected, provide additive contributions to the open-circuitvoltage. The layers of the optimal tandem cell (two-junction) havebandgaps at 1.3 um and 800 nm and take the limiting efficiency up to 44%compared to the 32% limiting efficiency of an optimal single-junctiondevice.

Quantum size-effect tuning provides one means of deploying materials ofa single composition to address the different needed spectral regions.The use of ternary and quaternary chalcopyrite nanoparticles adds afurther degree of freedom in harnessing efficiently the various band'swithin the sun's spectrum reaching the earth: the tuning ofstoichiometry. With bandgaps 740 nm and 1.23 eV, CuGaSe₂ and CuInSe₂ areexcellent candidates for the two junctions of a near-optimal tandemcell. Tuning CIGS stoichiometry along the continuum between CuGaSe₂ andCuInSe₂ provides continuous optimization over the intervening spectralrange.

The realization of the benefits described above is found through CGS,CIS, and CIGS colloidal nanoparticles exhi biting one of more of thefollowing characteristics: purity of phase (to generate a single bandgapwithin each junction); colloidal stability and a lack of aggregation (toensure practical processing); excellent crystallinity (to provide sharpabsorption onsets needed in optimized multijunction devices).

Compared with the well-developed synthesis of binary nanoparticles, theproduction of monodisperse ternary and quaternary chalcopyritenanoparticles in solution remains a challenge. CuGaSe₂ nanoparticleshave not previously been synthesized in the solution phase. Synthesis ofCuInSe₂ nanoparticles and nanorods in solution has been reported toproduce a wide dispersion in size. CIGS nanoparticles have beensynthesized in methanol or pyridine at low temperature and employed asprecursors for spray pryrolysis of CIGS thin film; however, thesematerials resulted in either large, uncontrolled aggregate formation,or, unless annealed at elevated temperatures, generated amorphousmaterials.

A description follows of the facile synthesis of phase-pure CuGaSe₂,CuInSe₂ and CIGS nanoparticles in oleyamine using hot injection methods.Monodisperse CuGaSe₂, and CIGS nanoparticles with good crystallinitywere synthesized for the first time. Low-cost commercial salts, seleniumpowder, and technical grade oleylamine were employed as precursors.

The strategy used in an embodiment relied on the careful choice ofprecursor and ligand combinations, and the success and failure ofvarious strategies confined the criticality of ensuring similarreactivities of the various components in successfully synthesizingmonodispersed and phase-pure ternaries and quaternaries. Without suchreactivity-matching, a plurality of phases and/or broad size andmorphology distribution of products are obtained. Specifically,monodisperse chalcopyrite nanoparticles were synthesized under thehypotheses that: there is rapid formation of chalcopyrite compounds inmixtures of appropriate precursors with liquid Se; Cu, In, and Ga saltsas well as selenium powders dissolve in oleyamine at elevatedtemperatures to enable chalcopyrite nanoparticle formation; theinjection ratio of precursors must favor of ternary and quaternarycompound formation and avoid binary compounds in a successful synthesis.

In a typical optimized synthesis of CIGS nanoparticles, 0.2 mmolCu(acac)₃ (acac-acetylacetate), 0.1 mmol In(acac)₃ and 0.1 mmolGa(acac)₃ were dissolved in 5 ml oleylamine at 80° C., producing a clearstock solution. Separately, 0.4 mmol selenium powder was added to 10 mloleylamine and heated to 250° C. until it turned clear orange. The stocksolution was swiftly injected into the selenium solution under vigorousstirring, and the heating mantle was turned off but left in place. Thesolution turned into black immediately, suggesting rapid nucleation. Thetemperature slowly dropped to 100° C. and then was heated back to 250°C. and incubated for 1 hour. The products were rinsed with methanol andtoluene twice, and finally dispersed in toluene.

The sizes, morphologies, crystallinities, and composition of thesynthesized particles were examined using transmission electronmicroscopy (TEM), selected area electron diffraction (SAED), and powderX-ray diffraction (XRD). FIG. 19A shows TEM images of chalcopyrite(CuGaSe₂) nanoparticies synthesized along with their corresponding SAEDpattern, under an embodiment. CuGaSe₂ nanoparticies were plate-like withirregular morphology and average size 11 nm. FIG. 19B shows TEM imagesof chalcopyrite (CuInSe₂) nanoparticles synthesized along with theircorresponding SAED pattern, under an embodiment. CuInSe₂ nanoparticleswere a mixture of triangular, deformed hexagonal, and round plate-likenanoparticles with average size 16 nm. FIG. 19C shows TEM images ofchalcopyrite (CIGS) nanoparticles synthesized along with theircorresponding SAED pattern, under an embodiment. CIGS nanoparticles hadaverage diameter 15 nm.

FIG. 21 shows plots of size distribution of as-synthesized nanoparticies(data based on manual counts of 80 nanoparticles from TEM images), underan embodiment. Each of these samples showed a narrow size distributionand their average sizes agree well with those calculated based on XRD.The spacing among nanoparticles in all samples was approximately 2.4 nm,twice the length of their passivating oleylamine ligand.

FIG. 20A shows powder XRD patterns of CuGaSe₂ 2004, CuInSe₂ 2002 andCIGS 2006 nanoparticles, under an embodiment; the vertical lines belowindicate the corresponding reflection peaks for bulkCuIn_(0.5)Ga_(0.5)Se₂ (JCPDS 40-1488), CuGaSe₂ (JCPDS 79-1809) andCuInSe₂ (JCPDS 40-1487). SAED (see the insets of FIG. 19) and XRD (seeFIG. 20A) confirmed excellent crystalline quality of all of theoptimally-synthesized particles of an embodiment. XRD diffraction peaksmatched very closely the bulk standard patterns of CuGaSe₂ (JCPDS79-1809), CuInSe₂ (JCPDS 40-1487) and Cu(In₀ ₅Ga_(0.5))Se₂ (JCPDS40-1488), respectively. In all cases, the match is with thelow-temperature tetragonal phase. FIG. 20B shows an ensemble ofUV-vis-NIR absorption spectrum of CuGaSe₂ 2004, CuInSe₂ 2002, and CIGS2006 nanoparticles in toluene, under an embodiment.

FIG. 22A shows TEM images of CuGaSe₂ synthesized by cooking Cu(Ac),Ga(acac)₃ and Se powder in oleylarnine at 250 C, under an embodiment.FIGS. 22B and 22C show TEM images of CuInSe₂ synthesized by cookingCu(Ac), In(Ac)₃ and Se powder in oleylamine at 250 C, under anembodiment. FIG. 23A shows TEM images and corresponding SAED of CuGaSe₂hexagonal microplates obtained in oleylamine and oleic acid mixture,under an embodiment.

FIG. 23B shows TEM images and corresponding SAED of CuInSe₂ hexagonalmicroplates obtained in oleylarnine and oleic acid mixture, under anembodiment. Figure S4A shows TEM images of CuInSe₂ nanoparticlessynthesized from Cu(acac)₂ and In(Ac)₃ precursors at 250° C. (scale barsare 50 nm), under an embodiment. FIG. 24B shows TEM images of CuInSe₂nanoparticles synthesized from Cu(Ac) and In(Ac)₃ precursors at 250° C.(scale bars are 50 nm), under an embodiment. FIG. 24C shows TEM imagesof CuInSe₂ nanoparticles synthesized from Cu(acac)₂ and In(acac)₂precursors at 250° C. (scale bars are 50 nm), under an embodiment. FIG.24D shows TEM images of CuInSe₂ nanoparticles synthesized from Cu(Ac)and In(acac)₃ precursors at 250° C. (scale bars are 50 nm), under anembodiment.

It was determined that the hot injection method, as well as the specificchoice of ligands and precursors described above, was crucial tosuccessful synthesis. When all precursors were mixed and heated, arelatively wide size dispersion, including large aggregates, resulted(Figure S2). Syntheses that employed ligands such as trioctylphosphine,trioctylphosphine oxide, dodecanethiol, oleic acid, stearic acid, andcombinations thereof, generally failed, with the exception thatsingle-crystalline CuInSe₂ nanoplates were obtained in a mixture ofoleylamine and oleic acid (see FIG. 23). Best results were obtainedusing Cu(acac)₂ and Ga(acac)₃ for CuGaSe₂. Best results were obtainedusing Cu(acac)₂ and InCl₃ for CuInSe₂. Best results were obtained usingCu(acac)₂, In(acac)₃ and Ga(acac)₃ for CIGS. Alternative precursorcombinations resulted in wide size distribution (seen by TEM) andmultiple phases of products (seen by comparison of XRD with standards)(see FIG. 24).

FIG. 25A shows TEM images of CIGS nanoparticles synthesized by injectionCu(acac)₂ In(acac)₃ and Ga(acac)₃ oleylamine solution into Se/oleylamineat an injection temperature of 270° C., under an embodiment. FIG. 25Bshows TEM images of CIGS nanoparticles synthesized by injectionCu(acac)₂, In(acac)₃ and Ga(acac)₃ oleylamine solution intoSe/oleylamine at an injection temperature of 220° C. under anembodiment. FIG. 26A shows TEM images of CIGS nanoparticies synthesizedwith a precursor ratio CIGE3228-0.15 mmol Cu(acac)₂, 0.1 mmol Ga(acac)₃and 0.1 mmol In(acac)₃ to 0.4 mmol Se, under an embodiment. FIG. 26Bshows TEM images of CIGS nanoparticles synthesized with a precursorratio CIGE4138-0.20 mmol Cu(acac)₂, 0.15 mmol Ga(acac)₃ and 0.5 mmolIn(acac)₃ to 0.40 mmol Se, under an embodiment. FIG. 26C shows XRDpatterns of CIGS nanoparticles synthesized with precursor ratiosCIGE3228-0.15 mmol Cu(acac)₂, 0.1 mmol Ga(acac)₃ and 0.1 mmol In(acac)₃to 0.4 mmol Se 2602 and CIGE4138-0.20 mmol Cu(acac)₂, 0.15 mmolGa(acac)₃ and 0.05 mmol In(acac)₃ to 0.40 mmol Se 2604, under anembodiment. Table 21 shows composition of CIGS nanoparticies of anembodiment calculated from Inductively Coupled Plasma Atomic EmissionSpectrometry (ICP).

In the embodiments described herein, temperature provided for tuning ofthe size and composition of the CIGS particles. Changing injectiontemperature between 220 and 270° C. tuned the average size ofas-synthesized CIGS from 12 to 18 nm (see FIG. 25). When the precursorratio of Cu:Ga:In:Se was changed from 4:2:2:8, 4:3:1:8 to 3:2:2:8 whilekeeping the injection temperature at 250° C., CuGa₀ ₃₈In₀ ₄₁Se_(1.87),CuGa₀ ₄₇In_(0.30)Se_(1.75) and CuGa_(0.41)In_(0.83)Se_(1.98) wereobtained, respectively (see FIG. 26 and FIG. 29). The observed deviationof product stoichiometry from precursor ratios is attributed to modestlydifferent reactivities of the metal precursors.

The absorption spectra of as-synthesized CuGaSe₂, CuInSe₂ andCu(InGa)Se₂ nanoparticles in toluene are presented in FIG. 20B, under anembodiment. The absorption of CuGaSe₂ and CuInSe₂ begin at 729 nm and1200 nm, consistent with their bulk bandgaps of 1.68 eV and 1.01 eV.Cu(InGa)Se₂ nanoparticles of an embodiment, however, showed strongabsorption in the visible and a weak tail in the near infrared possiblyattributable to free carrier absorption. FIG. 27 shows XRD patterns2702-2706 of CIGS nanoparticles arrested for different reactionduration, under an embodiment. The XRD characterization of productsobtained one (1) minute after injection in CIGS synthesis showedexclusively Cu(In_(0.5)Ga_(0.5))Se₂ diffraction peaks, implying thatchalcopyrite nanoparticles formed essentially immediately uponinjection.

Therefore, facile synthesis of monodispersed, crystalline, single-phaseCuGaSe₂, CuInSe₂ and CIGS nanoparticles was carried out by carefuladjustment of precursors and injection temperatures in oleylamine.Preliminary experiments indicate that the same approach may be appliedto CuInS₂ and CuGaS₂ via adoption of sulfur powder as the precursor.FIG. 28A shows representative TEM images and SAED pattern of CuInS₂nanoparticies produced in oleylamine using sulfur powder instead ofselenium powder, under an embodiment. FIG. 28B shows representative TEMimages and SAED pattern of CuGaS₂ nanoparticles produced in oleylamineusing sulfur powder instead of selenium powder, under an embodiment.Tuning of size and composition, combined with low cost of precursors andsimplicity of fabrication, are of interest in tandem photovoltaicdevices.

Experimental details of CuGaSe₂, CuInSe₂ and CIGS synthesis, isolationand characterization included the use of one or more of the followingchemicals, but the embodiment is not so limited: Copper(I) acetate(CuAc); copper (II) acetylacetate (Cu(acac)₂); indium (III) chloride(InCL₃); indium (III) acetylacetate (In(acac)₃); gallium (III)acetylacetate (Ga(acac)₃); selenium powder (Se); oleylamine (70%)(pumped under vacuum at 80° C. overnight before use).

Regarding preparation and isolation of CuGaSe2, CuInSe2 and CIGSnanoparticies, all synthesis was carried out with standard schlenk line.A round two-necked flask was located in a heating mantle, and one neckwas connected to a condenser while the other neck was sealed by septumand one thermocouple was used to control the temperature. For a typicalprocedure of CuGaSe₂ synthesis, 0.2 mmol Cu(acac)₂, 0.2 mmol Ga(acac)₃and 5 ml oleylamine were mixed at room temperature and kept at 80° C.under vacuum for 1 h to dissolve all the precursors completely. Thissolution was marked as solution A. Then 10 ml oleylamine and 0.4 mmol Sepowder were filled in a separate flask, and pumped at 120° C. for 0.5 hto further degas any residual air and/or moisture and then N₂ wasintroduced for the left whole reaction. The solution was heated up to250° C. in around 10 min and it gradually changed from colorless toorange to brownish red due to the dissolution of Se powder inoleylamine, which generally took about one hour. When no solid Se wasleft, 5 ml solution A in a syringe with 17 gauge needle was swiftlyinjected into the solution at 250° C. in about 5 second under vigorousstirring. The solution turned into black at once and turned off theheating mantle immediately (without removal of heating mantle).Temperature decreased slowly to 100° C. in around 15 min and then heatedthe solution back to 250° C. again and incubated for another 1 h. Thesolution was always black for the duration of the procedure. When thereaction was done, heating mantle was removed and the flask naturallycooled down to room temperature. Five (5) ml of anhydrous methanol wasinjected into the raw solution, and the solution was centrifugated at3000 rpm for 1 min. The clear supernantant was discarded and the blacksediments were redispersed in 10 ml anhydrous toluene. Five (5) mlmethanol was added to the solution again and isolated the nanoparticlesby centrifugation. The clear supernatant was decanted, the precipitateswere dispersed in 8 ml anhydrous toluene, and centrifugated at 3000 rpmfor 1 min and only the brown supernatant was collected.

As for the synthesis of other chalcopyrite nanoparticles, all theprocedures were kept exactly the same while 0.2 mmol InCl₃ for CuInSe₂,0.1 mmol In(acac)₃) and 0.1 mmol (Ga(acac)₃) for CuInGaSe₂ were used,respectively.

In performing structural, optical and electrical characterization ofchalcopyrite nanoparticles, transmission electronic microscopy (TEM) andselected area electron diffraction (SAED) were done on an EOL-2010-FEGmicroscope equipped with a tungsten filament and operated underaccelerating voltage of 200 kV. X-ray diffraction (XRD) was carried outin a Siemens diffractometer with Cu Kα radiation (λ=1.54175 Å).Vis-near-NIR absorption spectra of Chalcopyrite nanoparticles in toluenewere recorded at room temperature using a Cary 500 UV/vis/near-IRspectrophotometer. For ICP measurements, samples were acid digested,diluted with 18 mOhm water and assayed on a Perkin Elmer Model Optima3000DV ICP AEOS.

The embodiments described herein include a composite materialcomprising: semiconductor nanocrystals; and organic molecules thatpassivate the surfaces of the semiconductor nanocrystals, wherein atleast one property of the organic molecules facilitates the transfer ofcharge between the semiconductor nanocrystals.

The at least one property of an embodiment includes delocalization of atleast one type of charge carrier across at least a portion of theorganic molecules, wherein the at least one type of charge carrierincludes at least one of electrons and holes.

The composite material of an embodiment comprises at least one benzenering forming at least a portion of the organic molecules, wherein the atleast one benzene ring results in the delocalization of the at least onetype of charge carrier.

The semiconductor nanocrystals of an embodiment comprise at least one ofPbS, PbSe, PbTe, CdS, CdSe, CdTe, SnS, SnSe, SnTe, Si, GaAs, Bi2S3,Bi2Se3, CuInS₂, CuInSe₂, Cu(InGa)Se₂ (CIGS), CuGaSe₂.

The organic molecules of an embodiment comprise at least one ofBenzenedithiol, Dibenzenedithiol, Mercaptopropionic acid,Mercaptobenzoic acid, Pyridine, Pyrimidine, Pyrazine, Pyridazine,Dicarhoxybenzene, Benzenediamine, and Dibenzenediamine.

The embodiments described herein include a semiconductor material, thesemiconductor material comprising a p-type semiconductor materialincluding semiconductor nanocrystals, wherein at least one property ofthe semiconductor material results in a mobility of electrons in thesemiconductor material being greater than or equal to a mobility ofholes.

The embodiments described herein include a semiconductor material, thesemiconductor material comprising an n-type semiconductor materialincluding semiconductor nanocrystals, wherein at least one property ofthe semiconductor material results in a mobility of holes in thesemiconductor material being greater than or equal to a mobility ofelectrons.

The embodiments described herein include a device comprising asemiconductor material in contact with a first electrode and a secondelectrode, wherein the semiconductor material is a p-type semiconductormaterial comprising semiconductor nanocrystals, wherein properties ofthe semiconductor material result in a mobility of electrons in thesemiconductor material being greater than or equal to a mobility ofholes.

The embodiments described herein include a device comprising asemiconductor material in contact with a first electrode and a secondelectrode, wherein the semiconductor material is an n-type semiconductormaterial comprising semiconductor nanocrystals, wherein properties ofthe semiconductor material result in a mobility of holes in thesemiconductor material being greater than or equal to a mobility ofelectrons.

Unless the context clearly requires otherwise, throughout thedescription and the claims, the words “comprise,” “comprising,” and thelike are to be construed in an inclusive sense as opposed to anexclusive or exhaustive sense; that is to say, in a sense of “including,but not limited to.” Words using the singular or plural number alsoinclude the plural or singular number respectively. Additionally, thewords “herein,” “hereunder,” ‘“above,” “below,” and words of similarimport, when used in this application, refer to this application as awhole and not to any particular portions of this application. When theword “or” is used in reference to a list of two or more items, that wordcovers all of the following interpretations of the word: any of theitems in the list, all of the items in the list and any combination ofthe items in the list.

The above description of embodiments is not intended to be exhaustive orto limit the systems and methods to the precise forms disclosed. Whilespecific embodiments of, and examples for, the embodiments are describedherein for illustrative purposes, various equivalent modifications arepossible within the scope of the systems and methods, as those skilledin the relevant art will recognize. The teachings of the embodimentsprovided herein can be applied to other systems and methods, not onlyfor the systems and methods described above.

The elements and acts of the various embodiments described above can becombined to provide further embodiments. These and other changes can bemade to the embodiments in light of the above detailed description.

In general, in the following claims, the terms used should not beconstrued to limit the embodiments to the specific embodiments disclosedin the specification and the claims, but should be construed to includeall systems that operate under the claims. Accordingly, the embodimentsare not limited by the disclosure, but instead the scope of theembodiments is to be determined entirely by the claims.

While certain aspects of the embodiments are presented below in certainclaim forms, the inventors contemplate the various aspects of theembodiments in any number of claim forms. Accordingly, the inventorsreserve the right to add additional claims after filing the applicationto pursue such additional claim forms for other aspects of theembodiments.

What is claimed is:
 1. A composite material, comprising: at least onenanoparticle; organic molecules to passivate surfaces of the at leastone nanoparticle, at least one property of the organic molecules beingto facilitate a transfer of charge between ones of the at least onenanoparticle, each of the organic molecules being a benzene ring thatincludes at least one functional group; and at least one bidentatelinker coupled to the nanoparticle and consisting of a first end groupand a second end group, the bidentate linker including a conjugatedmoiety coupled between the first end group and the second end group. 2.The composite material of claim 1, wherein the at least one nanoparticleis chosen from at least one of the molecules including lead sulfide(PbS) and lead selenide (PbSe).
 3. The composite material of claim 1,wherein the at least one bidentate linker is a dithiol.
 4. The compositematerial of claim 1, wherein the at least one bidentate linker is abenzenedithiol.
 5. The composite material of claim 1, wherein electronicproperties of the composite material are stable in time.
 6. A device,comprising: a first electrode; a composite material including at leastone nanoparticle; and at least one bidentate linker having a first endgroup and a second end group, the bidentate linker including aconjugated moiety coupled between the first end group and the second endgroup; a second electrode forming a metal-semiconductor junction with alayer that includes colloidal quantum dots; and a benzene ring thatincludes at least one functional group.
 7. The device of claim 6,wherein the first electrode is chosen from at least one of the groupsincluding indium tin oxide (ITO), platinum (Pt), and gold (Au).
 8. Thedevice of claim 6, wherein the second electrode is chosen from at leastone of the groups including aluminum (Al), silver (Ag), and manganese(Mg).
 9. The device of claim 6, wherein an external quantum efficiencyof the device exceeds about 40%.